Current aircraft braking systems utilize either steel or carbon disks that serve as the friction materials and heat sinks. steel-based systems were the original brake materials. They were used on all aircraft until the emergence of carbon-carbon (C—C) composite materials in the 1970s. C—C composites are now the state-of-the-art material for aircraft brake heat sinks and are being used in the vast majority of new military and large commercial aircraft programs.
Ceramic Matrix Composites (CMCs) exhibit some extraordinary thermal and mechanical properties and hold the promise of being outstanding materials for aircraft brake friction applications, as well as attractive candidates for the next-generation heat-sink materials for such applications. A particular CMC system that indeed possesses the potential for use as a next-generation aircraft brake material, and offers potential breakthrough performance, has recently been identified. In particular, brake materials based on a slurry cast/melt-infiltrated boron carbide matrix composite system have been shown to offer extremely attractive benefits relative to both steel and carbon brake materials.
As a class of materials, ceramics are known to possess low density, high hardness and high oxidation resistance; some of them also have attractive heat capacity and thermal conductivity. Compared to the carbon-carbon composites used today, ceramics have the potential of providing some key performance advantages in terms of reduced wear rate, enhanced oxidation resistance, and reduced heat-sink weight and/or volume.
The earliest attempts to use using ceramics for aircraft braking system applications were based on monolithic ceramics and cermets. However, none of these attempts were successful. The major cause for the resultant failures was due to the inadequate mechanical properties, especially low impact resistance and low fracture toughness, associated with the well-known characteristic brittleness of ceramics. Thus, activities on the next-generation heat-sink materials for aircraft braking system applications have focused on the development of fiber-reinforced CMCs that would improve the fracture toughness and impact resistance (reducing the brittleness) while retaining the other advantages of ceramics.
The two prime candidate CMC material systems, identified for aircraft braking system heat-sink applications, are respectively based on silicon carbide (SiC) and boron carbide (B4C) due to thermo-mechanical considerations. Among these two material systems, the B4C-based CMCs have the particular attractions that B4C is the third hardest material known, with only diamond and cubic boron nitride being harder, and that it has a heat capacity greater than both SiC and carbon. However, B4C-based fiber-reinforced CMCs suitable for aircraft brake application had heretofore never been made due to processing difficulties associated with B4C. Any previous attempts known were limited to materials without fiber reinforcements. For example, U.S. Pat. No. 5,878,849 issued Mar. 9, 1999, discussed infra., describes a cermet material made by infiltrating a pressed preform of B4C powder (not filament or fiber) with aluminum; the end product is proposed for aircraft brake uses.
Silicon-filled CMCs have been reported in both U.S. and foreign literature to show improved friction coefficients and/or wear life in certain configurations. See, for example: R. W. Froberg and B. A. Grider, “High Friction carbon/carbon Aircraft Brakes”, 40th Int. SAMPE Symp., May 8-11, 1995, extended abstracts, pp 942-944; R. W. Froberg and T. E. Pratt, “Brake System with Improved Brake Material”, U.S. Pat. No. 4,815,572 issued Jul. 24, 1987 (assigned to Parker Hannifin Corp); W. Krenkel, “CMC Materials for High Performance Brakes”, ISTA Conference on Supercars, Aachen, 31 Oct.-4 Nov. 1994 (paper from the German Aerospace Research Establishment Institute of Structures and Design), Stuttgart; A. Lacombe, “Friction System Using Refractory Composite Material”, U.S. Pat. No. 5,007,508 issued Apr. 15, 1991 (assigned to SEP, France).
For example, a helicopter brake with a higher static friction coefficient (μ) than C—C, and a stable μ of 0.30 at all energies including a “rejected take-off” (RTO), was reported by Parker Hannifin Corp. (See R. W. Froberg and B. A. Grider, ibid., and U.S. Pat. No. 4,815,572, ibid.) Chopped carbon fibers were molded, carbonized, and densified to 1.60-1.65 g/cm3 by carbon vapor infiltration (CVI), and then reacted with liquid Si at 1850° C. to form SiC to a depth of 0.06-0.07 inch.
In 1994, the German Aerospace Research Establishment reported sub-scale dynamometer results on a C—C+SiC composite which showed improved stability, lower wear, and shorter processing times than C—C. (See W. Krenkel, ibid.) Pyrolyzed resin-impregnated carbon fiber preforms, infiltrated with Si at 1500° C., yielded composites containing −35% SiC by weight. Friction coefficients varied between 0.2 and 1.0, higher than for C—C under comparable conditions, increasing with decreasing velocity. Wear was not affected by temperature up to 900° C.
Lower net wear rates were disclosed by a patent to SEP covering aircraft brakes in which a C—C composite disk is worn against a disk containing C or Sic fibers and the CVI matrix consists of SiC as the principal phase with minor amounts of C or BN on the fibers. (See U.S. Pat. No. 5,007,508 ibid.)
However, while the foregoing examples illustrate the potential advantages of Si-based CMC's, few of these claims have been independently substantiated. Very often, the friction and wear (F&W) test duty cycle, including load, pressure, or length of testing time and number of cycles were either not reported, or were far less severe than those demanded under realistic aircraft braking conditions. Furthermore, many of these studies only cited either the friction or the wear results, by themselves, instead of the more relevant combined F&W data.
Options in processing boron carbide have been reviewed by Thevenot. See: F. Thevenot, “Boron Carbide—A Comprehensive Review”, 1989, pp 2.1-2.23; F. Thevenot, “Formation of Carbon-Boron Bonds”, in Inorganic Reactions and Methods, ed. J. Zuckerman, A. Hagen, VCH Publishers, New York, 1989, 10, pp 2-11; F. Thevenot, “Sintering of Boron Carbide and Boron Carbide—Silicon Carbide Two-Phase Materials and their Properties”, J. Nucl. Mater., 1988, 152, pp 154-162. High thermal conductivity and strength require minimum porosity and attention to the B/C ratio in the solid solution. Monolithic B4C has often been hot pressed from fine powder in graphite dies at 2375-2475° C., which would not permit inclusion of fiber reinforcements without damage. Hot isostatic pressing (HIP) with glass encapsulated molds of Ti was used to densify B4C with excess C at 1700° C. (See: H. T. Larker, L. Hermansson and J. Adlerborn, “Hot Isostatic Pressing and its Applicability to Silicon Carbide and Boron Carbide”, Mater. Sci. Monoar., 1987 38A pp 795-803; ASEA. Final densities >99% were achieved by sintering without encapsulation at 2000° C. under 200 MPa (29 ksi) Argon pressure for 2 hours. (See K. A. Schwetz, W. Greliner, and A. Lipp, “Mechanical Properties of HIP Treated Sintered Boron Carbide”, Inst. Phys. Conf., 1986, Series No. 75, Chap. 5, pp 413-425.) Pressureless sintering with additives that form lower melting borides still required high temperatures, and exaggerated grain growth led to low strength. Furthermore, similar to hot pressing, pressureless sintering is not suitable for the inclusion of fiber reinforcements due to volume shrinkage.
Carbon fiber-reinforced B4C for Tokamak fusion reactors was prepared at Los Alamos in 1978-1979 by hot pressing multiple layers of coated low-modulus graphite cloth (UCC WCA) at 2100° C. and under a pressure of 32 MPa (4.6 ksi). See: L. R. Newkirk, et. al., “Preparation of Fiber Reinforced Titanium Diboride and Boron Carbide Composite Bodies”, Proc. 7th CVD Con., 1979, pp 515-521. [Proc. 7th Eur Con. CVD, 1989]; R. E. Riley, et. al, “Preparation and Uses of Amorphous Boron Carbide Coated Substrates”, U.S. Pat. No. 4,287,259 issued Sep. 1, 1981. Hot pressed billets with up to 43 vol % B4C and 1.87 g/cm3 density were obtained; however, fibers were damaged. A 20-cloth laminate with 37 vol % B4C, hot pressed at 2050° C. and under a pressure of 4.6 ksi for 15 min, had flexural strengths of only 7.4-9.9 ksi.
There are several general reviews of boron carbide-carbon vapor deposition (CVD) studies, which have been concerned primarily with coating surfaces at higher temperatures. See: H. Hannache et. al., “Kinetics of Boron Carbide Chemical Vapor Deposition and Infiltration”, Proc. 5th European Conf. on CVD, 1985, pp 219-233; A. W. Moore and H. F. Volk, “Chemical Vapor Deposition of Boron Carbide”, AMMRC CR 69-10, August, 1969; L. C. Vandenbulcke, “Theoretical and Experimental Studies on the Chemical Vapor Deposition of Boron Carbide”, Ind. Eng. Chem. Prod. Res. Dev., 1985 24, pp 568-575; U. Jansson, “Chemical Vapor Deposition of Boron Carbides”, Materials & Manufacturing Processes, 6(3), 1991, pp 481-500. Of relevance to infiltrated preforms is information on the effects of deposition conditions and gas ratios on the structure, hardness, and other properties, which vary with the B/C ratio in the deposits. CVI requires sufficient supersaturation to avoid depletion, and a low deposition rate giving a process limited by the surface reaction rate. (See H. Hannache et. al., ibid.) So far, none of the processes cited were useful for densifying fiber or filament preforms.
In summary, prior to the work resulting in the present invention, as described hereinafter in more detail, there was no known practical processing technique for producing dense, fiber-reinforced B4C-based CMCs. It appears that, as a result, no relevant F&W data has been published for B4C-based systems. The development of a suitable fabrication process for fiber-reinforced B4C CMCs is an important challenge to be overcome.
Additions to carbon fibers of submicron B4C and ZrO2 powders, and SiC from pre-ceramic polymers, liquid Si, SiO vapor and CVI SiC, plus Si—O—C sol-gel, were evaluated all under internally funded R&D programs. Initially, overall results were mixed. Some promising results were achieved with the CVI-SiC additions, but reproducible levels of significant F&W improvement could not be obtained. In addition, there were concerns about the anticipated high cost of SiC CVI.
The addition of ceramics to carbon fiber preforms from silicon-based liquid precursors was investigated as a lower cost alternative to SiC CVI. Formation of a silicon-oxy-carbide (SiOC) via sol-gel processing resulted in lower wear than materials made by other liquid precursor methods. A heat sink was made with SiOC, and was densified by CVI-C. Strength was acceptable, and dynamometer tests showed 20% lower average wear rate after wear-in than commercially available Boeing 747 C—C brake materials. However, this approach had the drawback of producing a material with low thermal conductivity. Heat treatment to improve conductivity degraded material strength and wear rate.
A cermet is a mixture of ceramic and metal powders, usually very fine powders approaching sub-micron grain size, which are co-sintered, at elevated temperatures, usually in a liquid phase. U.S. Pat. No. 5,878,849 issued Mar. 9, 1999 to Prunier, Jr. et al., assigned to The Dow Chemical Company, discloses boron carbide-aluminum cermets and a process for producing such. In the development of the disclosed process, Al+B4C particle mixtures were slip cast and heated to 1400° C. to form Al—B, Al—C, and Al—B—C phases with some residual Al solid solution. For other related cermet development, see: D. C. Halverson, A. J. Pyzik, and I. A. Aksay, “Processing and Microstructural Characterization of B4C—Al Cermets”, Ceram. Eng. Sci. Proc., 1985, 6, pp 736-744; D.C. Halverson, A. J. Pyzik and I. A. Aksay, “Boron-Carbide-Aluminum and Boron-Carbide-Reactive Metal Cermets”, U.S. Pat. No. 4,605,440 issued Aug. 12, 1986; A. J, Pyzik and I. A. Aksay, “Multipurpose Boron Carbide-Aluminum Composite and its Manufacture via the Control of the Microstructure”, U.S. Pat. No. 4,702,770 issued Oct. 28, 1987. A rapid omni-directional compaction process for cermets was also studied. See: A. J. Pyzik and A. Pechenik, “Rapid Omnidirectional Compaction of Ceramic-Metal Composites”, Ceram. Eng. Sci. Proc., 9, (7-8) pp 965-974 [1988]; S. Ashley, “Ceramic-Metal Composites: Bulletproof Strength”, Mechanical Engineering, July, 1990, pp 46-51.
A recent attempt was made to reproduce the material publicly disclosed in U.S. Pat. No. 5,878,849, ibid.; the material was fabricated as described in that patent, and subjected to industry standard tests. To wit, the material was slip cast, and “green” bodies were machined to shape before Al infiltration and heat treatment. Final dimensioning of the very hard material was by electrical discharge machining and diamond grinding.
The heat capacity of the cermets of U.S. Pat. No. 5,878,849 was 47% higher than that of C at 700° F. and 5% higher at 1040° F., although thermal conductivity was 20% lower. Tests of this cermet on the HSFT apparatus, described infra., indicated a relatively stable friction coefficient (μ) with wear rates ⅓-½ that of carbon. Nevertheless, low frequency vibration was encountered and most samples fractured during test. When run against itself on a full-scale friction screening machine (FSM), μ was in the range required for aircraft brakes. However, the test disks were found to cut grooves into one another at stress concentration points, and structural failure later followed. Furthermore, when tested against a C—C disk, the initial μ was 0.37 for taxi stops and 0.29 for service stops; but after −300 stops, a drastic decline of μ to unacceptably low levels was encountered. The reason for this severe performance degradation was not clear. A full-scale (PC-12) hybrid brake, with a ceramic rotor tested vs. C—C stators, failed by cracking at the outer radius of the rotor friction face after only one cycle.
The prominent features of the material disclosed and taught by U.S. Pat. No. 5,878,849 can be summarized as follows:    No fiber reinforcements is used.    The article is characterized by a non-continuous metal phase. Aluminum alloys are the most common metal phase, but silicon is also cited.    The article comprises a continuous ceramic phase. Boron carbide in the Dow system is evaluated, as stated above, but other ceramic systems are also cited.    No CVD material is present.    Solid state sintering of the ceramic phase is critical for the formation of a continuous ceramic phase.    The articles exhibit fracture toughness −5 MPa m1/2    The ceramic phase is 85 to 98 vol % of material.    The diameter of the metal regions is only 0.25 to 30 micrometers.
Certain CMC systems are disclosed in a series of U.S. Patents issued to Singh et al, assigned to The General Electric Company. Those patents all disclosed the use inter alia of carbon fiber preforms. As disclosed in those patents, in some embodiments boron carbide can be used. These General Electric patents fall generally into one of two groups. The first group all produce solid state sintered ceramic bodies, wherein the composite matrix is densified by hot-press sintering, and wherein the final sintered body is reduced in size from the body before sintering. Those patents, all of which disclose technology which is unsuitable for producing complex, near-net shape ceramic matrix components, are: U.S. Pat. No. 4,886,682, “Process for Producing a Filament-Containing Composite in a Ceramic Matrix”, issued Dec. 12, 1989; U.S. Pat. No. 4,915,760, “Method of Producing a Coated Fiber-Containing Composite”, issued Apr. 10, 1990; U.S. Pat. No. 4,931,311, “Method of Obtaining a Filament-Containing Composite with a Boron Nitride Coated Matrix”, issued Jun. 5, 1990; U.S. Pat. No. 5,051,301, “Coated Fiber-Containing Composite”, issued Sep. 24, 1991; U.S. Pat. No. 5,067,998, “Fibrous Material-Containing Composite”, issued Nov. 11, 1991; U.S. Pat. No. 5,160,676, “Fibrous Material-Containing Composite”, issued Nov. 3, 1992; and U.S. Pat. No. 5,407,734, “Fiber-Containing Composite”, Apr. 18, 1995.
The second group of General Electric patents, as referred to above, discloses the infiltration of a porous body with a molten silicon infiltrant. As disclosed, a limitation is imposed in all cases on the fibrous materials that can be utilized. That limitations is that all must be first coated entirely with boron nitride so as to avoid a reaction and bonding between the silicon infiltrant and the fibrous material, that reaction and bonding which would, according to the disclosure, destroy fiber pull-out capabilities and thus destroy fracture toughness. Those patents are U.S. Pat. No. 4,889,686, “Composite Containing Coated Fibrous Material”, issued Dec. 26, 1989; U.S. Pat. No. 4,944,904, “Method of Obtaining a Fiber-Containing Composite”, issued Jul. 31, 1990; U.S. Pat. No. 4,981,822, “Composite Containing Coated Fibrous Material”, issued Jan. 1, 1991; U.S. Pat. No. 5,021,367, “Fiber-Containing Composite”, issued Jun. 4, 1991; U.S. Pat. No. 5,043,303, “Filament-Containing Composite”, issued Aug. 27, 1991; U.S. Pat. No. 5,330,854, Filament-Containing Composite”, issued Jul. 19, 1994; U.S. Pat. No. 5,376,427, “Ceramic Composite Containing Coated Fibrous Material”, issued Dec. 27, 1994; U.S. Pat. No. 5,387,299, Ceramic Composite Containing Coated Fibrous Material”, issued Feb. 7, 1995; and U.S. Pat. No. 5,432,253, “Composite Containing Fibrous Material”, issued Jul. 11, 1995.
The fabrication and evaluation of SiC-based CMC materials was investigated. The fabrication process used was a variation of a slurry casting/melt infiltration (SC/MI) process previously developed by The Carborundum Company for the preparation of SiC/SiC composites. See: U.S. Pat. No. 5,296,311, “Silicon Carbide Reinforced Reaction Bonded Silicon Carbide Composite”, McMurtry et al., issued Mar. 22, 1994; U.S. Pat. No. 5,436,042, “Shaped Green Ceramic Fabric Preform Segments for Fiber Reinforced Composite Article”, Lau et al., issued Jul. 25, 1995; U.S. Pat. No. 5,484,655, “Aluminum Nitride-Coated Silicon Carbide Fiber”, Lau et al., issued Jan. 16, 1996; U.S. Pat. No. 5,643,514, “Process for Manufacturing a Silicon Carbide Composition”, Chwastiak et al., issued Jul. 1, 1997; U.S. Pat. No. 5,817,432, “Silicon Carbide Reinforced Reaction Bonded Silicon Carbide Composite”, Chwastiak et al., issued Oct. 6, 1998; U.S. Pat. No. 5,840,221, “Process for Making Silicon Carbide Reinforced Silicon Carbide Composite”, Lau et al., issued Nov. 24, 1998; and U.S. Pat. No. 5,945,062, Silicon Carbide Reinforced Reaction Bonded Silicon Carbide Composite”, Chwastiak et al., Aug. 31, 1999.
The Carborundum CMC system starts with a SiC fiber-reinforced preform coated with a CVD BN, AlN or TiB2 interface coating. This preform is then impregnated with an aqueous slurry containing Sic powders with a bi-modal particle size distribution. The slurry-impregnated preform is then heated to −1410° C. and infiltrated with molten silicon. With a hold time of 30 minutes or less, the final infiltration results in a near full-density CMC, with a SiC—Si two-phase matrix, commonly referred to as a Melt-Infiltrated SiC ceramic matrix composite (MI/Sic CMC).
Specimens were prepared by the present inventors in accord with the foregoing published Carborundum process, however the Carborundum process was modified to include carbon fiber preforms rather than the Sic fiber preforms taught by that process. Further, the carbon fiber preforms were CVI coated with carbon, rather than being coated with BN, AlN or TiB2 as taught by the Carborundum process.
The specimens were tested with the MI-SiC CMC specimens wearing against one another, the MI-SiC CMC specimens wearing against carbon-carbon specimens, and the MI-SiC CMC specimens wearing against a hybrid C—C/CVI Sic material. In general, all of the tests indicated that the MI-SiC CMC material had high and stable friction coefficients but imparted much higher wear rates than C—C. All testing was conducted on the High Speed Friction Tester (HSFT) using conditions commonly used by The BFGoodrich Company to evaluate C—C friction materials, as described infra. in the next paragraph.
In the development of the fiber-reinforced B4C CMCs of the present invention, and for testing of the SiC-based MI-CMC specimens prepared by using the foregoing Carborundum-like process, a High Speed Friction Tester (HSFT) was used to perform friction and wear (F&W) screening on rotor/stator pairs of 0.375-inch thick samples with a 1.55 in2 friction interface, 2.25 inches OD and 1.75 inches ID. The disks are mounted on ceramic insulators, and a thermocouple in the rotor, 0.05 inch from the wear face, records temperature, normally 500-1500° F. Normal force and torque on the stator are measured in a series of stops from 5000 rpm (43.6 ft/sec), which last ˜50 seconds. Data are recorded each 0.008 sec. and averaged each 0.55 sec.; these averages are stored, and the average friction coefficient μ, the rms deviation from average μ, the temperatures and the loads were recorded for a series of stops at fixed pressure, normally between 38 psi and 76 psi on the friction surface. Loads up to 130 psi were applied to obtain high temperatures representative of normal energy or RTO stop conditions. Thickness change at 7 locations around the wear face is measured after each series with a micrometer. A typical standard “up” test sequence used for evaluating CMC materials is 400 stops at 13 psi, 200 stops at 23 psi, 200 stops at 35 psi, 200 stops at 47 psi, 200 stops at 58 psi, 200 stops at 69 psi, and (if an acceptable amount of the wear surface remains and oxidation of the material is limited) 5 stops at 130 psi (-RTO condition). The use of the term “up” indicates that the pressure is increased for each set of stops, beginning at 13 psi and ending up at 69 psi (or 130 psi if acceptable wear surface remains). Typical “down” testing for evaluating CMC materials would be 200 stops at 69 psi, followed by 200 stops at 58 psi., followed by 200 stops at 47 psi, followed by 200 stops at 35 psi, followed by 200 stops at 23 psi and then followed by 400 stops at 13 psi. Typical “up” and “down” testing are depicted on FIGS. 1 and 2 of the drawings. Average wear rate (mils per surface per 1000 stops) is compared for each series, and wear rates after wear-in are compared, usually for the last 200 stops. The sub-scale, high-speed friction tester (HSFT) was used to assess the effects of ceramic additions on friction coefficient (μ) and wear of carbon brakes.
As mentioned above, this series of F&W tests on the HSFT was completed on the SiC-based MI-CMC specimens prepared by using the foregoing Carborundum-like process. Also, as mentioned above, all the test results showed that the MI-SiC CMC materials had high and stable friction coefficients. However, while low wear rates were achieved at low brake pressures, the wear rates at higher pressures often jumped to orders of magnitude higher than that of a typical C—C material.
Post-test characterization revealed that large Sic particles found in the MI-SiC CMC matrix were responsible for the high wear rate. It was curious to find that although the starting SiC powders used in the slurry were very fine (mostly below one micron in size, with a small fraction in the 5 micron range), large SiC crystals (significantly larger than 20 microns) were prevalent in the CMC matrix after processing (see photo micrograph in FIG. 3). On the other hand, most of the starting fine SiC powder particles had virtually disappeared. The presence of the large (≧20 microns) SiC crystals, instead of the <1 micron SiC particles, may be causing the increase in the wear rate. Since these large SiC crystals were more abrasive than the finer particles, an analogous behavior would be found in the use of coarse grit sand paper instead of fine grit sand paper.
Detailed examination of the post-melt-infiltration material indicated the occurrence of the C interaction with liquid Si, and recrystallization, a process resulting in a microstructure containing large (>20 micron) alpha-Sic particles embedded in large “pools” of Si.
The fact, that these particles were much larger (>20 micron) than the original alpha-Sic particles (mostly <1 micron) used in the starting SiC slurry, is indicative of an interaction between liquid Si and the CVD carbon coating during the melt infiltration process followed by the process of recrystallization of the Sic.
Experimental work has indicated that the following mechanism may be occurring in the application of the modified Carborundum process, as outlined above:    (1) The CVD carbon material deposited on the carbon preform may be reacting with liquid silicon to form a solid SiC layer at the C/Si interface. After that, any further reaction would have to depend on the diffusion of either C or Si reactant through this solid SiC product layer so formed. This would be a very slow process at the low temperatures used in the MI processing (−1410° C.).    (2) Simultaneously, as the thickness of the SiC layer is increasing, some of the SiC formed may also be in the process of being dissolved in liquid silicon at the liquid Si/solid Sic interface (interface 1). However, since liquid silicon has a very low solubility for SiC (less than a few hundred ppm at 1410° C.), such a dissolution process should normally stop very quickly once the saturation limit is reached.    (3) On the other hand, because the slurry contains many alpha SiC particles to start with, they may now be acting as sinks or “seeds” for the growth of larger SiC crystals. The dissolved SiC (hereinafter referred to as SiC) might thus be precipitating on these seeds, leading to the growth of even larger SiC crystals.    (4) With the dissolved SiC precipitating out from the liquid silicon and recrystallizing onto the SiC seeds, the SiC concentration near the crystalline liquid Si/solid SiC crystal interface (interface 2) may become lower; as a result, a SiC concentration gradient may be established in the Si melt between interfaces 1 and 2. This may enable a sustained transport of SiC down the concentration gradient.
The net result may be a reaction, dissolution, and recrystallization cycle that might lead to the growth of large SiC crystals. Also, due to thermodynamic considerations, such a recrystallization and growth process may be favored to occur on the surfaces of the larger SiC particles, and most of the small SiC particles from the original SiC slurry may also be dissolved via a similar process. The end result may be a matrix with a prevalence of large, recrystallized SiC crystals and little or no fine SiC particles.
There is a need to be able to reduce the ceramic particle size, in the MI-Ceramic CMC matrix, to reduce the high wear rate at high brake pressures to a wear rate that is comparable to that of typical C—C materials, if not better.